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Article

TiAl-Based Oxidation-Resistant Hard Coatings with Different Al Contents Obtained by Vacuum-Pulse-Arc Granule Melting

by
Alexander N. Sheveyko
,
Konstantin A. Kuptsov
,
Philipp V. Kiryukhantsev-Korneev
*,
Maria N. Fatykhova
,
Georgy M. Markov
and
Dmitry V. Shtansky
Scientific Educational Center of SHS, National University of Science and Technology “MISIS”, Leninsky Prospect 4/1, 119049 Moscow, Russia
*
Author to whom correspondence should be addressed.
Coatings 2024, 14(1), 6; https://doi.org/10.3390/coatings14010006
Submission received: 1 December 2023 / Revised: 12 December 2023 / Accepted: 16 December 2023 / Published: 19 December 2023
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Abstract

:
A method was proposed for increasing the oxidation resistance of promising wrought Ti2AlNb ortho-alloys by depositing γ-TiAl-based coatings. Using original vacuum pulse-arc melting of 100 μm thick granule layers, coatings with different Al/Ti ratios and a thickness of 50–60 µm were obtained on the surface of the Ti50Al25Nb25 alloy. Granules Ti50Al44Nb4.9Mo1B0.1 (at.%), 20–60 μm in size, were employed. To vary Al content, initial granules and their mixture with Al powder were used. Excellent adhesion of the coatings is ensured by the similar chemical composition and structure of the substrate and coatings, as well as micro-metallurgical reactions between granules and the substrate that occur during treatment. The resulting coatings had a submicron gradient structure consisting of TiAl and Ti3Al intermetallic compounds. During oxidation at 850 °C for 10 h, an oxide layer consisting of a mixture of α-Al2O3, TiO2, and AlNbO4 was formed on the coating surfaces. With an increase in the annealing duration to 100 h, a dense α-Al2O3 oxide layer, approximately 0.5 µm thick, was formed over the primary oxide mixture, the quality of which was higher in coatings enriched with aluminum.

1. Introduction

High-temperature alloys based on TiAl and Ti3Al intermetallic compounds are of high interest for the manufacture of turbine blades and aircraft engines due to the combination of low density (3.7–4.5 g/cm3) and high oxidation resistance [1,2,3]. However, a major drawback of these alloys is their poor machinability and formability [4,5,6,7].
Currently, the most promising and technologically advanced alloys are based on the orthorhombic intermetallic Ti2AlNb compound [8,9,10]. However, these alloys exhibit lower oxidation resistance compared to γ-TiAl alloys, since their oxidation is accompanied by the formation of Nb oxides (AlNbO4 и Nb2O5) [11,12]. Intermetallic γ-TiAl-based alloys, due to higher Al content, provide long-term oxidation resistance at temperatures up to 850 °C [13]; however, their application in the manufacturing of turbine blades is limited due to the formation of fatigue cracks [14]. Moderate alloying of γ-TiAl material with Nb (4%–6%) and Mo (1%–2%) (so-called TNM alloys) opens up opportunities for improving their plastic deformation [15] but reduces oxidation resistance [16,17]. Alloying of these alloys with micro-additives of various elements is aimed at increasing mechanical properties and oxidation resistance. For example, micro-additions of boron ensure the formation of TsiB2 particles, which leads to a grain refinement and, as a consequence, to an increase in high-temperature creep resistance [18,19]. In addition to boron, Y2O3 micro-additions are capable of suppressing grain-boundary oxygen diffusion and stabilizing α-Al2O3 [6]. The positive effect of introducing multiple elements (W, B, Y, and C, Si) on the microstructure and oxidation behavior of Nb-containing TiAl alloys have been demonstrated [20]. It is noteworthy that attempts to enhance the oxidation resistance of TiAl-based alloys are mainly carried out with an aluminum content < 50%. According to the Ti-Al phase diagram, an increase in Al content above 53% transfers the alloy into the two-phase region with the formation of TiAl3 and TiAl2 intermetallics, in which a significant change in their ratio is observed with temperature variations, which has a detrimental effect on the fatigue strength of oxidation-resistant alloys.
An increase in the Al content in the surface layer without deteriorating the bulk material properties can be achieved by forming Al-rich coatings. The deposition of TiAl-based oxidation-resistant coatings can be realized by various methods; for example, hot-dip aluminizing [21] or diffusion batch cementation method [22]. Both methods provide high adhesion but lead to damage of the carefully engineered alloy structure. Selective Ti etching with chlorine and fluorine, aimed at increasing the Al content in the surface layer, can be accompanied by leaching of substrate elements and deteriorate mechanical properties [23]. A similar technology of selective Ti removal from surface layers by sulfurization in an H2S–H2 atmosphere leads to the formation of a sulfide layer which requires an additional step for its removal [24,25].
Physical vapor deposition (PVD) is the most versatile technology for depositing oxidation-resistant Al-rich coatings onto TNM alloys. A wide range of different oxidation-resistant PVD coatings has been obtained [13,26,27]. However, the thickness of these coatings does not exceed 10 μm, which significantly limits their use for long-term high-temperature applications.
At elevated temperatures, interdiffusion of elements between the substrate and coating can lead to significant structural and composition changes resulting in deterioration of oxidation resistance. For this reason, coatings with an elemental composition close to that of the substrate are more preferable; for example, PVD TiAl coatings on TNM alloys [28,29]. In addition, the Al amount in thin coatings required to ensure the formation of aluminum oxide is limited during long-term operation [30]. Therefore, a coating with perfect adhesion to the surface, increased Al content, and a relatively large thickness of 50–100 μm is required.
Electrospark deposition (ESD) is one of the coating methods that meet these requirements well. For example, Al-rich coatings on the surface of Ni superalloys have been obtained by treating both Al [31] and intermetallic NiAl electrodes [32]. An important feature of ESD is discharge melting and mixing of the substrate and consumable electrode materials. This ensures ideal coating adhesion and partially heals substrate defects, which is critical for additive methods. High process purity and increased productivity are achieved when ESD is conducted in a vacuum [33,34]. An increase in the ESD coating thickness can be achieved by increasing the discharge power. However, treatment intensification is accompanied by an increase in surface roughness, which imposes a thickness limit of a high-quality coating with an acceptable level of surface roughness, usually about 50 μm. The disadvantage of this method is the active mechanical mixing of the melt containing products from both the electrode and the substrate. This interaction leads to the incorporation of substrate melt fragments into the upper coating layers. Intensive mixing of the melt can be eliminated by non-contact melting of a pre-applied powder layer by pulses with energy typical of an ESD. This technology has been successfully validated to deposit a NiAl layer to protect a Ni-based superalloy from oxidation [35].
The originality of this work lies in the development of a non-contact pulsed melting method for the formation of Al-rich coatings on the Ti-Al-Nb alloy. This method is unique in that it avoids the use of traditional arc melting techniques, which can lead to undesirable defects and inhomogeneities in the coating. The use of granules containing Nb, Mo, and B ensures that the composition of the deposited layer is well controlled, while the mixing of pure Al powder with the granules allows for the intentional increase in Al content in the coating. This approach is particularly beneficial for applications where high Al content is required for improved oxidation resistance or mechanical properties.

2. Materials and Methods

Oxidation-resistant Ti2AlNb ortho-alloy with a composition of Ti50Al24Nb25V1 (at.%) was utilized as a substrate. Samples with dimensions of 60 × 20 × 5 mm3 were cut from cast ingots and then were subjected to a heat treatment to stabilize the ortho-phase. Afterwards, one side of the substrate was ground using abrasive paper to a roughness (Ra) of 6 µm. Using the method of arc plasma spheroidization of powders obtained by self-propagating high-temperature synthesis (SHS), granules with composition Ti50Al44Nb4.9Mo1B0.1 (at.%) and a size of 20–60 µm were obtained [36]. Oxygen contamination of the initial SHS powders was at a level of 0.6–0.8 at.%, while after spheroidization in Ar + H2 plasma it decreased to 0.2–0.3 at.%.
To vary Al content in the coatings, either only granules or a mixture of granules with Al powder (size 10–16 μm, 5 at.% residual O) processed in a ball mill were used, which resulted in approximately 50 at.% of Al content in the coating. Granules or a mixture of granules and Al powder were deposited on a horizontal substrate surface using thin calibrated spacers on the sample sides to ensure a constant layer thickness of 100 μm. For fixing on the substrate, the granule layer was annealed in a vacuum furnace at a temperature of 600 °C for 1 h at a residual pressure of 10−3 Pa.
The coatings were deposited using an experimental vacuum unit equipped with a 3-axis CNC module, which allows scanning the surface along a given trajectory with a non-consumable electrode [35]. A tungsten electrode with a diameter of 2 mm was located at a constant distance (discharge gap) of 300 ± 50 µm from the granule layer. Original discharge electrical circuit was used (Figure 1a). To initiate a breakdown of the vacuum gap between the electrode and the granules, high-voltage pulses were used, followed by a powerful discharge of the energy stored in the capacitors into a channel pre-formed by the breakdown. Induction coils were included in the circuit to increase the discharge time and limit the peak current. The current pulse shape is illustrated in Figure 1b. The vacuum chamber was first evacuated to a pressure of 1 × 10−3 Pa, and then Ar was supplied to a pressure of 1 × 104 Pa (0.1 atm). Relatively high pressure made it possible to achieve better focusing of the arc discharge. Melting of the granule layer was carried out under the following parameters: pulse frequency of 5 Hz, voltage of 400 V, pulse energy of 15 J, electrode speed of 1 mm/s, and a scanning step of 0.2 mm.
The structure and elemental composition of the coatings were examined by X-ray diffraction (XRD) using a D8 diffractometer (Bruker, Billerica, MA, USA) in Bragg–Brentano geometry with Cu-Kα radiation, transmission electron microscopy (TEM) using a JEM 2100 instrument (JEOL, Tokyo, Japan), scanning electron microscopy (SEM) and energy-dispersive X-ray (EDX) spectroscopy (EDS) on a JSM 7600F electron microscope (JEOL, Tokyo, Japan) equipped with an EDX detector. The content of light elements, such as B, C, and O, was additionally determined by glow discharge optical emission spectrometry (GDOES) using a Profiler-2 optical spectrometer (Horiba, Kyoto, Japan). GDOES was also employed to analyze the elemental distribution throughout the coating thickness. To study the coatings using transmission electron microscopy (TEM), samples were prepared using a focused ion beam (FIB) using a FEI Scios Dualbeam scanning electron microscope (FEI, Hillsboro, OR, USA).
Hardness (H) was measured in a Testing Laboratory for Functional Surfaces (MISIS) using a Nano Hardness Tester (CSM Instruments, Peuseux, Switzerland) equipped with a Berkovich indenter under a load of 10 mN. The hardness distribution over the coating thickness was measured in cross-sections from the surface to the substrate along a line at an angle of 45° with a distance of 8 µm between prints (13 measurements in a row).
To evaluate the oxidation resistance, the coatings were annealed at 850 °C for 10 and 100 h in a SNOL 7.2/1200 furnace (SnolTherm, Utena, Lithuania). For TEM studies, the samples were first coated with a carbon mask to protect the upper oxide layer and then placed in a FIB setup to prepare the foil.

3. Results and Discussion

3.1. Microstructure of Coating Obtained by Melting Granules

SEM images of the granule layer on the surface before treatment are shown in Figure 2a. The size of spherical granules varies from 12 to 55 μm. Inside the powder mixture, an uneven distribution of Al powder is observed in the form of agglomerates up to 50 μm in size (Figure 2b).
Figure 2 shows SEM cross-section and top view images of the coating surfaces. Both coatings exhibit similar surface morphology (fish scale type), consisting of solidified melt areas approximately 300 µm in size caused by individual discharge pulses. On the coating surface, individual incompletely melted granules are also observed, which are not found inside the coating. The coatings have a thickness of approximately 50 µm. Individual cracks are observed in the cross-sections, some of which extend to the surface. It is noteworthy that these cracks do not propagate deep into the substrate. The observed cracks can negatively affect the oxidation resistance. The coating can be divided into an upper dark layer and a lighter layer adjacent to the substrate. Since no clearly defined interface between the coating and the substrate was found, it can be assumed that during the contact of the granule melt with the substrate, a thin substrate layer also melts. The composition of various zones inside the coatings (Table 1) confirms the active interaction of the granule melt with the substrate, as evidenced by the increase in Nb content in the transition zone compared to the granule composition. The coating has a gradient structure, in which the bottom coating layer contains a significantly higher amount of Ti and Nb compared to the surface layer. The composition of the top coating layer is identical to the composition of the original granules/mixtures. The minimum Ti/Al ratio (maximum Al content) is achieved in a coating additionally alloyed with Al powder. The structure of the top layer in both coatings consists of submicron grains darker surrounded by a network of phases with a darker contrast (inserts).
The granule-melting diagram (Figure 2f) illustrates the stages of the process. The primary high-voltage discharge breaks down the gap along the shortest distance from the electrode to the nearest granule. When the voltage during the initiating breakdown drops below the capacitor voltage, diode D3 (Figure 1a) opens, directing all electric energy stored in the capacitors into a pre-formed discharge channel. An inductive coil limits the discharge current, preventing boiling and splashing of the melt. During the pulse, the arc discharge is shifted to areas with lower resistance, resulting in both direct heating of the granules and their indirect heating and melting through heat transfer from the previously molten layer (depicted by semicircular arrows).
The results of XRD analysis of the coatings are presented in Figure 3a. The two main phases observed in the coatings are intermetallic compounds γ-TiAl (ICDD 65-0428) and α2-Ti3Al (ICDD 14-0451). In addition to these phases, small amounts of a β-Ti-based solid solution (ICDD 89-4913) and a phase with an fcc lattice identified as (Ti, Nb) (ICDD 88-2330) [37,38] are observed. The substrate contains two phases: the orthorhombic phase Ti2AlNb (COD 96-152-2559) and a β-Ti-based solid solution with high Nb content.
GDOES profiles illustrate, in depth, element distributions in the direction from the surface to the substrate (Figure 3b). A monotonic change in the element concentrations is observed throughout the entire coating thickness. The composition of the upper coating layers correlates with EDS data. A substrate-matched composition is achieved at an etch depth of 55 µm, which corresponds to the observed coating thickness on polished samples.
To explain the phase formation in TiAl and TiAl + Al coatings, the isothermal section of the Ti-Al-Nb ternary phase diagram at 700 °C and the section of the ternary phase diagram at a constant Nb content of 5 at.% were used (Figure 4).
Comparison of EDS and XRD data with phase diagrams allows a more complete explanation of the coating structure. It is important to consider that during the ESD process, the melt zones are rapidly cooled due to heat transfer into the substrate, which leads to a rapid solidification. Upon cooling, coatings with a lower Al content fall into the two-phase region of coexistence of two intermetallic compounds γ-TiAl and α2-Ti3Al, the ratio of which in the coating depends on the cooling rate according to the phase diagram, only γ-TiAl is expected to form in a coating with a higher Al content. However, according to XRD data, in addition to the main γ-TiAl component, Ti3Al phase is also present in the coating. This is most likely due to the fact that the area of XRD analysis includes the lower coating layers located closer to the substrate, which contain a significantly higher Ti content (Table 2).

3.2. Annealing and Oxidation

To study temperature-activated structural changes in the coatings and the formation of oxide layers on the surface, the coatings were annealed in air at 850 °C for 10 and 100 h. SEM images of the cross-sections of annealed coatings with different Al content are presented in Figure 5, and the results of the EDS composition analysis in selected regions are shown in Table 2. After annealing, a surface oxide layer with a thickness of approximately 1.5 µm is observed (Figure 5c).
In the lower part of both coatings, precipitates of a new phase are observed, which have a brighter contrast and a higher Nb content (Figure 5d, Table 2, Point 1). The main structural elements have gray (Figure 5d, Table 2, point 3) and dark (Figure 5d, Table 2, point 2) contrast and correspond to Ti3Al and TiAl, respectively.
Figure 6 shows indenter print images and the dependence of microhardness on the depth of the analyzed layer in TiAl and TiAl + Al coatings before and after annealing at 850 °C for 100 h. Hardness measurements were carried out starting from a depth of 6 μm. The initial coatings have a hardness of about 10 GPa, which is twice the hardness of the substrate. A monotonous increase in hardness from the substrate to the coating surface is observed both in the initial and annealed states, which is explained by a gradual increase in the volume fraction of the TiAl phase. After annealing, the hardness of both coatings increases to 13–14 GPa, which is associated with upward Al diffusion, stabilization of intermetallic phases, and an increase in TiAl fraction (Figure 6d). The increased hardness in the lower part of the coatings may be associated with dispersion strengthening due to the precipitation of the Nb-rich phase.
Optical images of the TiAl + Al coating before and after annealing in comparison with oxidized substrate are shown in Figure 7, coating SEM cross-sectional images, demonstrating the microstructure of the oxide layers are depicted in Figure 8, and their XRD patterns are presented in Figure 9. During annealing, a loose oxide layer, approximately 6 µm thick, is formed on the surface of an uncoated substrate, containing titanium oxides, rutile and anatase, as well as a small of the AlNbO4 phase. This oxide layer has weak adhesion to the substrate and can be easily removed by wiping with a cloth (as shown by the circle in Figure 7). In this layer, longitudinal delaminations and nanosized pores are observed (Figure 8a,b).
Table 2. Content of main elements in different parts of the TiAl + Al coating and oxide layers.
Table 2. Content of main elements in different parts of the TiAl + Al coating and oxide layers.
at.%FigureTiAlNbO
TiAl + Al, pt. 1Figure 5463519
TiAl + Al, pt. 2Figure 552408
TiAl + Al, pt. 3Figure 557349
Oxide on substr., pt. 4Figure 8198965
Oxide TiAl + Al, pt. 5 greyFigure 82025352
Oxide TiAl + Al, pt. 6 whiteFigure 82014660
Oxide TiAl + Al, pt. 7 blackFigure 81128556
The surface oxide layers after annealing, regardless of annealing time (10 or 100 h), have a thickness of ~1.5–2 µm and consist of several layers, the composition and structure of which depend on the coating composition and the oxidation time (Figure 8c–f). After 10 h of oxidation, two types of grains with bright and dark contrast are observed in the structure of both coatings. The Ti content in both components is approximately 20 at.%, while the Al content in dark grains is twice as high as in light grains, 25 and 14 at.%, respectively (Table 2, Points 5, 6). The grain size in the near-surface layer does not exceed 50 nm (Figure 8c,e), and in the lower oxide layer it is 100–300 nm.
According to XRD analysis, the oxide layers on all the coatings contain two main components: TiO2-rutile (ICDD 76-0319) and Al2O3-corundum (ICDD 05-0712). In the region of small 2θ angles, low intensity peaks from the AlNbO4 phase (ICDD 51-0023) can be seen (Figure 9b). After long-term annealing (100 h), a double-layer oxide is formed on the coating surface (Figure 8d,f). The lower oxide layer also consists of a mixture of TiO2 and Al2O3 phases, while the upper homogeneous layer with a thickness of 0.5 μm is mainly formed by aluminum oxide. We note that in the oxide layer of the coating with a lower Al content, nanosized inclusions of other oxides are observed.
To study in more detail the microstructure of the oxide layers in the Al-enriched coating after annealing for 100 h (Figure 10), TEM studies were carried out. Selected area electron diffraction (SAED) patterns and EDS data were acquired from four different areas: (I) the top oxide layer (Figure 10f), (II) TiO2 and (III) Al2O3 grains in a mixed oxide layer (Figure 10g,h), and (IV) the non-oxidized coating area located beneath the oxide layer (Figure 10i). Identification of these areas was performed using EDS elemental mapping (Figure 10b–d) and element content in marked areas (Figure 10e).
The top single-phase oxide layer, about 300 nm thick, consists of fine α-Al2O3 grains with a hexagonal R3c crystal structure. No other structural elements were found in this layer. Below is a zone with a thickness of 1.2 μm, consisting of a mixture of α-Al2O3 and TiO2 with a rutile crystal structure with the size of individual oxide grains being 300 nm. The coating layer (IV) located beneath the oxide corresponds to the Ti3Al intermetallic compound.
The formation of a thin surface layer of alumina can be explained as follows. The oxide layer is mostly based on a mixed oxide of titanium and aluminum, formed due to the close values of Gibbs free energy (−780 and −850 kJ/mol, respectively) [39,40]. This oxide does not act as an effective barrier to oxygen diffusion [41,42]. The oxidation resistance of TiAl intermetallics can be improved by alloying with Zr, Nb, Mo, Hf, Ta, W, and Re [17].
At the initial moment of oxidation, a mixed oxide of TiO2 and Al2O3 is formed. As the oxide thickness increases, oxygen atoms penetrate the mixed oxide and, as a result of counter-diffusion, react with the more active Al, forming Al2O3 within the oxide layer.
The introduction of Nb into TiAl reduces the probability of Al2O3 formation compared to TiO2 [43]. As a result, the formation of Al2O3 within the oxide layer is suppressed and Al atoms have the opportunity to reach the surface and form a dense surface oxide.

3.3. Final Remarks

The possibility of a new technology for melting TiAl-based granules for depositing protective coatings to increase the oxidation resistance of Ti2AlNb ortho-alloy was demonstrated. The utilization of multielement granules instead of a mixture of metal powders makes it possible to reduce the oxygen content in the coating (<1 at.%). The Al content in coatings can be adjusted by adding some Al powder. The key advantage of the proposed technology is its implementation in a vacuum, which additionally ensures low oxygen and nitrogen contents in the coatings. An important feature of the pulsed arc melting technology is a non-contact simultaneous melting of granules and the substrate and the absence of intensive mixing of the melt. Vacuum deposition with cathodic polarity of the granule layer and the substrate ensures excellent spreading of the melt over the surface of the substrate [34]. It is important that the composition of the top coating layer predominantly corresponds to the initial composition of the granules.
The similarity of the chemical and phase composition of the coating and the substrate leads to the absence of a sharp interface and minimal boundary stresses due to the similarity of thermal expansion coefficients and elastic moduli, which provides excellent adhesive strength. Excellent oxidation resistance is achieved by using a mixture of TiAl-based granules and Al powder, which provides a high Al content. The formation of the upper oxide layer leads to Al-depletion in the underlying layer. In addition, long-term exposure at 850 °C reduces the Al content close to the coating/substrate interface [30], which can degrade the mechanical properties. Thus, the undoubted advantage of thick coatings over thin ones is that they have sufficient reserves of Al, which can stabilize Al-based phases in the coating over a long period under cross-diffusion of elements.
The Al content on the coating surface was 40 at.% (TiAl) and 49 at.% (TiAl + Al), which provided an Al/Ti ratio of 0.75 and 1.14, respectively. After oxidation in air at 850 °C for 100 h, an 1.5 μm thick oxide layer formed on the surface of both coatings, consisting of a thin (~0.5 μm) top layer of Al2O3 and a mixture of Al2O3 (corundum) and TiO2 (rutile) beneath. Annealing experiments also showed that the thickness of the dense oxide layer formed after 10 h does not change with time (up to 100 h), which confirms the excellent oxidation resistance of TiAl and TiAl + Al coatings at 850 °C.

4. Conclusions

A new technology for the deposition of oxidation resistant coatings was proposed based on vacuum pulse melting of Ti50Al44Nb4.9Mo1B0.1 (at.%) granules with a size of 20–60 μm with the addition of Al powder, which ensured the formation of thick (~50 μm) and hard (~10 GPa) coatings with a high aluminum content and low oxygen and nitrogen content. A 100 μm thick layer of granules or a granules/powder mixture was melted in vacuum by arc discharge pulses while scanning the surface with a tungsten electrode. High adhesion strength of coatings to the substrate was achieved due to the simultaneous melting of granules and a thin substrate layer, as well as the similarity of their elemental composition. After oxidation at 850 °C for 10 h, a two-phase Al2O3 + TiO2 layer with a thickness of 1.5 μm was formed on the surface. The thickness of the oxide layer remained unchanged after annealing for 100 h, but it transformed into a dense upper Al2O3 layer and a lower layer of a Al2O3 + TiO2 mixture. The large coating thickness, which provided reserves of Al, allowed us the possibility to count on long-term preservation of its microstructure during long-term operation at high temperatures under conditions of inevitable element diffusion.

Author Contributions

Conceptualization, Funding acquisition, A.N.S.; Investigation, A.N.S., K.A.K., P.V.K.-K., M.N.F. and G.M.M.; Visualization, K.A.K. and P.V.K.-K., Writing—original draft, A.N.S. and K.A.K.; Review and editing, D.V.S. and K.A.K. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the Russian Science Foundation, grant number No. 22-29-00757.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Data are contained within the article.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. Electrical diagram of the vacuum-pulse-arc discharge circuit (a) and shape of the current pulse (b).
Figure 1. Electrical diagram of the vacuum-pulse-arc discharge circuit (a) and shape of the current pulse (b).
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Figure 2. SEM images of a granule layer before (a) and after mixing with Al powder (b). SEM cross-sectional images of TiAl (c) and TiAl + Al (d) coatings, surface morphology of TiAl (e) and TiAl + Al (f) coatings, as well as scheme of melting process (g).
Figure 2. SEM images of a granule layer before (a) and after mixing with Al powder (b). SEM cross-sectional images of TiAl (c) and TiAl + Al (d) coatings, surface morphology of TiAl (e) and TiAl + Al (f) coatings, as well as scheme of melting process (g).
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Figure 3. XRD spectra (a) and GDOES profiles (b) of as-deposited TiAl and TiAl + Al coatings.
Figure 3. XRD spectra (a) and GDOES profiles (b) of as-deposited TiAl and TiAl + Al coatings.
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Figure 4. Isothermal section of the Ti-Al-Nb ternary phase diagram at 700 °C (a) and a fragment of the section of the Ti-Al-Nb phase diagram at a constant Nb content of 5 at.% (b) [30]. Red allows and yellow/red circles show the composition of coatings.
Figure 4. Isothermal section of the Ti-Al-Nb ternary phase diagram at 700 °C (a) and a fragment of the section of the Ti-Al-Nb phase diagram at a constant Nb content of 5 at.% (b) [30]. Red allows and yellow/red circles show the composition of coatings.
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Figure 5. SEM cross-sectional images of TiAl (a) and TiAl + Al (b–d) coatings after annealing in air at 850 °C for 100 h, enlarged images of the oxide layer on the surface (c) and the coating microstructure at the boundary of the transition zone (d).
Figure 5. SEM cross-sectional images of TiAl (a) and TiAl + Al (b–d) coatings after annealing in air at 850 °C for 100 h, enlarged images of the oxide layer on the surface (c) and the coating microstructure at the boundary of the transition zone (d).
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Figure 6. Distribution of microhardness over the coatings thickness (a), images of indents on cross-sections of TiAl (b) and TiA + Al (c) coatings, and XRD patterns of coatings (d) before and after annealing at 850 °C for 100 h.
Figure 6. Distribution of microhardness over the coatings thickness (a), images of indents on cross-sections of TiAl (b) and TiA + Al (c) coatings, and XRD patterns of coatings (d) before and after annealing at 850 °C for 100 h.
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Figure 7. Optical image of TiAl + Al coating before annealing, and the substrate and TiAl + Al coating after annealing at 850 °C for 100 h.
Figure 7. Optical image of TiAl + Al coating before annealing, and the substrate and TiAl + Al coating after annealing at 850 °C for 100 h.
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Figure 8. SEM cross-sectional images of the substrate (a,b), TiAl (c,d) and TiAl + Al (e,f) coatings after annealing at 850 °C for 100 h (a,b,d,f) and 10 h (c,e).
Figure 8. SEM cross-sectional images of the substrate (a,b), TiAl (c,d) and TiAl + Al (e,f) coatings after annealing at 850 °C for 100 h (a,b,d,f) and 10 h (c,e).
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Figure 9. XRD patterns of annealed TiAl and TiAl + Ti coatings: 10–90° 2θ (a) and on an enlarged scale at 10–37.5° 2θ (b).
Figure 9. XRD patterns of annealed TiAl and TiAl + Ti coatings: 10–90° 2θ (a) and on an enlarged scale at 10–37.5° 2θ (b).
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Figure 10. Cross-sectional TEM images (a) showing the microstructure of TiAl + Al coating after annealed at 850 °C for 100 h, corresponding EDS elemental mapping (b–d), Ti, Al, and O content in marked areas (e), and corresponding SAED patterns (f–i).
Figure 10. Cross-sectional TEM images (a) showing the microstructure of TiAl + Al coating after annealed at 850 °C for 100 h, corresponding EDS elemental mapping (b–d), Ti, Al, and O content in marked areas (e), and corresponding SAED patterns (f–i).
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Table 1. Elemental composition of substrate, granule and coating as determined by EDS and GDOES.
Table 1. Elemental composition of substrate, granule and coating as determined by EDS and GDOES.
at.%TiAlNbMoB *Fe *C *O *Ti/Al
Substrate502525∑ < 12
Granule50444.91.00.1∑ < 11.14
Coating TiAl surface53395.00.90.10.20.81.01.36
Coating TiAl trans. zone5329160.90.10.10.40.51.83
Coating TiAl + Al surface43495.00.60.10.31.01.00.88
Coating TiAl +Al trans. zone523511.50.40.10.10.50.41.49
* GDOES data.
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MDPI and ACS Style

Sheveyko, A.N.; Kuptsov, K.A.; Kiryukhantsev-Korneev, P.V.; Fatykhova, M.N.; Markov, G.M.; Shtansky, D.V. TiAl-Based Oxidation-Resistant Hard Coatings with Different Al Contents Obtained by Vacuum-Pulse-Arc Granule Melting. Coatings 2024, 14, 6. https://doi.org/10.3390/coatings14010006

AMA Style

Sheveyko AN, Kuptsov KA, Kiryukhantsev-Korneev PV, Fatykhova MN, Markov GM, Shtansky DV. TiAl-Based Oxidation-Resistant Hard Coatings with Different Al Contents Obtained by Vacuum-Pulse-Arc Granule Melting. Coatings. 2024; 14(1):6. https://doi.org/10.3390/coatings14010006

Chicago/Turabian Style

Sheveyko, Alexander N., Konstantin A. Kuptsov, Philipp V. Kiryukhantsev-Korneev, Maria N. Fatykhova, Georgy M. Markov, and Dmitry V. Shtansky. 2024. "TiAl-Based Oxidation-Resistant Hard Coatings with Different Al Contents Obtained by Vacuum-Pulse-Arc Granule Melting" Coatings 14, no. 1: 6. https://doi.org/10.3390/coatings14010006

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